EP4029959A1

EP4029959A1 - High-strength steel sheet, and method for producing same

Info

Publication number: EP4029959A1
Application number: EP20880186.0A
Authority: EP
Inventors: Kazuki ENDOH; Yoshiyasu Kawasaki; Yuki Toji
Original assignee: JFE Steel Corp
Current assignee: JFE Steel Corp
Priority date: 2019-10-23
Filing date: 2020-10-08
Publication date: 2022-07-20
Also published as: KR20220066364A; JPWO2021079754A1; WO2021079754A1; CN114555845B; CN114555845A; MX2022004669A; EP4029959A4; US20220364196A1; JP7168073B2

Abstract

An object is to provide a high strength steel sheet having a tensile strength of 980 MPa or more and excellent formability and a method for manufacturing the high strength steel sheet.Provided is a high strength steel sheet that has a predetermined chemical composition and is manufactured under optimum conditions, the high strength steel sheet having a steel microstructure including, by area, ferrite: 30% or more and 80% or less, tempered martensite: 3.0% or more and 35% or less, and retained austenite: 8% or more, wherein the quotient of the area fraction of grains of the retained austenite, the grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less, divided by the total area fraction of the retained austenite is 0.3 or more, wherein the quotient of the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite is 1.5 or more, and the product of the quotient of the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite and the average aspect ratio of the retained austenite is 3.0 or more, wherein the quotient of the average C content (mass%) in the retained austenite divided by the average C content (mass%) in the ferrite is 3.0 or more, wherein the quotient of the average C content (mass%) in the retained austenite divided by the average Mn content (mass%) in the retained austenite is 0.05 or more, and wherein the content of diffusible hydrogen in steel is 0.3 ppm by mass or less.

Description

Technical Field

The present invention relates to a high strength steel sheet having excellent formability which is suitably used as a member in various industries, such as the automotive industry and the electrical industry, and a method for manufacturing the high strength steel sheet. Specifically, the present invention attempts to produce a high strength steel sheet that has a TS (tensile strength) of 980 MPa or more and is excellent in terms of not only ductility but also stretch flangeability and bendability.

Background Art

Improving fuel economy of automobiles has been an important issue from the viewpoint of global environmental conservation. Accordingly, there have been active movements to increase the strengths of car body materials in order to allow reductions in the thicknesses of the car body materials and the weights of car bodies. Since increasing the strength of a steel sheet may degrade the formability of the steel sheet, the development of a material having both high strength and high formability has been anticipated.
There has been proposed a high strength steel sheet that uses strain induced transformation of retained austenite, which has a high strength and excellent ductility. This steel sheet has a microstructure including retained austenite and can be readily formed into an intended shape due to retained austenite during forming, while the steel sheet comes to have a high strength subsequent to forming as a result of transformation of retained austenite to martensite.
For example, Patent Literature 1 proposes a high strength steel sheet having a tensile strength of 1000 MPa or more, a total elongation (EL) of 30% or more, and markedly high ductility, the steel sheet being manufactured using strain induced transformation of retained austenite. Such a steel sheet is manufactured by causing a steel sheet containing C, Si, and Mn as fundamental constituents to form austenite, subsequently quenching the steel sheet in the bainite transformation temperature range, and performing isothermal holding, that is, an austempering treatment. Retained austenite is formed as a result of C being concentrated at austenite due to the austempering treatment, and the formation of a large amount of retained austenite requires addition of a large amount of C exceeding 0.3%. However, an increase in the C concentration in steel degrades spot weldability, and in particular a C concentration exceeding 0.3% significantly degrades spot weldability. Thus, it has been difficult to bring this technique into active use for automotive steel sheets. In addition, no consideration is given to stretch flangeability and bendability in Patent Literature 1 because the main purpose is to improve the ductility of the high strength steel sheet.
In Patent Literature 2, a high Mn steel is used and a suitable strength-ductility balance is achieved by performing a heat treatment in the ferrite-austenite dual phase region. However, in Patent Literature 2, no study is performed on the improvement of ductility by concentrating Mn at untransformed austenite. Thus, there is a room for improving workability.
In Patent Literature 3, a medium Mn steel is used and total elongation is increased by performing a heat treatment in the ferrite-austenite dual phase region, concentrating Mn at untransformed austenite, and thereby forming stable retained austenite. However, no consideration is given to achievement of compatibility of elongation, stretch flangeability, and bendability. Moreover, no study is performed on the improvement of stretch flangeability and bendability by controlling the distribution of not only Mn but also C in the second phase composed of retained austenite and martensite. It is considered that the concentration of Mn is not at the level enough to achieve the compatibility of elongation, stretch flangeability, and bendability because the amount of time during which the heat treatment is performed in the manufacturing method described in Patent Literature 3 is small and the diffusion rate of Mn is low.
Furthermore, in Patent Literature 4, a medium Mn steel is used and uniform elongation and stretch flangeability are improved by performing a heat treatment of a hot rolled steel sheet in the ferrite-austenite dual phase region for a long period of time, facilitating the concentration of Mn at untransformed austenite, and thereby forming retained austenite grains having a high aspect ratio. However, in Patent Literature 4, the improvement of the ductility and stretch flangeability of a high strength steel sheet only by Mn concentration is studied; no study is performed on the compatibility of the improvement of stretch flangeability, bendability, and elongation by controlling the distribution of C and Mn in the second phase composed of retained austenite and martensite.

Citation List

Patent Literature

PTL 1: Japanese Unexamined Patent Application Publication No. S61-157625
PTL 2: Japanese Unexamined Patent Application Publication No. H1-259120
PTL 3: Japanese Unexamined Patent Application Publication No. 2003-138345
PTL 4: Japanese Patent No. 6123966

Summary of Invention

Technical Problem

The present invention was made in light of the above-described circumstances. An object of the present invention is to provide a high strength steel sheet having a TS (tensile strength) of 980 MPa or more and excellent formability and a method for manufacturing the high strength steel sheet. Note that the term "formability" used herein refers to ductility, stretch flangeability, and bendability. Solution to Problem
In order to achieve the above object, the inventors of the present invention conducted extensive studies in terms of the chemical composition of the steel sheet and a method for manufacturing the steel sheet for manufacturing a high strength steel sheet having excellent formability and consequently found the following facts.
Specifically, the inventors of the present invention found that it is important to heat a steel slab including Mn at a content of 2.50% by mass or more and 8.00% by mass or less and optional other alloying elements, such as Ti, at adequate contents, hot roll the steel slab at a finish rolling delivery temperature of 750°C or more and 1000°C or less, then perform coiling at 300°C or more and 750°C or less, optionally perform holding at a temperature equal to or lower than the Ac₁ transformation temperature for more than 1800 s, then perform cold rolling, subsequently perform holding at a temperature equal to or higher than the Ac₃ transformation temperature - 50°C for 20 s or more and 1800 s or less, then perform cooling to a cooling stop temperature equal to or lower than the martensitic transformation start temperature, then perform reheating to a reheating temperature of 120°C or more and 450°C or less, subsequently perform holding at the reheating temperature for 2 s or more and 600 s or less, and then perform cooling to room temperature such that film-like austenite grains at which C is concentrated, which serve as cores of fine retained austenite grains having a high aspect ratio in the subsequent annealing step, are formed. After the cooling has been performed, holding is performed to a temperature equal to or higher than the Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less, cooling is performed to 100°C or less, holding is then performed at more than 100°C and 400°C or less for 10 s or more, and then cooling is performed. It was found that the above-described method enables the production of a high strength steel sheet having excellent formability, the steel sheet having a steel microstructure including, by area, ferrite: 30% or more and 80% or less, tempered martensite: 3.0% or more and 35% or less, and retained austenite: 8% or more, wherein the quotient of the area fraction of grains of the retained austenite, the grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less, divided by the total area fraction of the retained austenite is 0.3 or more, wherein the quotient of the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite is 1.5 or more, and the product of the quotient of the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite and the average aspect ratio of the retained austenite is 3.0 or more, wherein the quotient of the average C content (mass%) in the retained austenite divided by the average C content (mass%) in the ferrite is 3.0 or more, wherein the quotient of the average C content (mass%) in the retained austenite divided by the average Mn content (mass%) in the retained austenite is 0.05 or more, and wherein the content of diffusible hydrogen in steel is 0.3 ppm by mass or less.
The present invention was made on the basis of the above-described findings. The summary of the present invention is as follows.

[1] A high strength steel sheet including a chemical composition containing, by mass, C: 0.030% or more and 0.250% or less, Si: 0.01% or more and 3.00% or less, Mn: 2.50% or more and 8.00% or less, P: 0.001% or more and 0.100% or less, S: 0.0001% or more and 0.0200% or less, N: 0.0005% or more and 0.0100% or less, and Al: 0.001% or more and 2.000% or less, with the balance being Fe and incidental impurities; a steel microstructure including, by area, ferrite: 30% or more and 80% or less, tempered martensite: 3.0% or more and 35% or less, and retained austenite: 8% or more, wherein a quotient of an area fraction of grains of the retained austenite, the grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less, divided by a total area fraction of the retained austenite is 0.3 or more, wherein a quotient of an average Mn content (mass%) in the retained austenite divided by an average Mn content (mass%) in the ferrite is 1.5 or more, and a product of the quotient of the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite and an average aspect ratio of the retained austenite is 3.0 or more, wherein a quotient of an average C content (mass%) in the retained austenite divided by an average C content (mass%) in the ferrite is 3.0 or more, wherein a quotient of the average C content (mass%) in the retained austenite divided by the average Mn content (mass%) in the retained austenite is 0.05 or more, and wherein a content of diffusible hydrogen in steel is 0.3 ppm by mass or less.
[2] The high strength steel sheet described in [1], wherein the chemical composition further contains at least one element selected from, by mass, Ti: 0.200%, Nb: 0.200% or less, V: 0.500% or less, W: 0.500% or less, B: 0.0050% or less, Ni: 1.000% or less, Cr: 1.000% or less, Mo: 1.000% or less, Cu: 1.000% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ta: 0.100% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, and REM: 0.0050% or less.
[3] The high strength steel sheet described in [1] or [2], the high strength steel sheet further including a galvanized layer disposed on a surface of the high strength steel sheet.
[4] The high strength steel sheet described in [3], wherein the galvanized layer is a galvannealed layer.
[5] A method for manufacturing a high strength steel sheet, the method including heating a steel slab having the chemical composition described in [1] or [2], hot rolling the steel slab at a finish rolling delivery temperature of 750°C or more and 1000°C or less, then performing coiling at 300°C or more and 750°C or less, subsequently performing cold rolling, then performing holding at a temperature equal to or higher than an Ac₃ transformation temperature - 50°C for 20 s or more and 1800 s or less, subsequently performing cooling to a cooling stop temperature equal to or lower than a martensitic transformation start temperature, then performing reheating to a reheating temperature of 120°C or more and 450°C or less, subsequently performing holding at the reheating temperature for 2 s or more and 600 s or less, then performing cooling to room temperature, subsequently performing holding at a temperature equal to or lower than an Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less, then performing cooling to 100°C or less, then performing holding at higher than 100°C and 400°C or less for 10 s or more, and subsequently performing cooling.
[6] The method for manufacturing a high strength steel sheet described in [5], the method further including, subsequent to the coiling and prior to the cold rolling, performing holding at a temperature equal to or lower than the Ac₁ transformation temperature for more than 1800 s.
[7] The method for manufacturing a high strength steel sheet described in [5] or [6], the method further including, subsequent to the holding at the temperature equal to or higher than the Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less, performing cooling, then performing a galvanizing treatment, and subsequently performing the cooling to 100°C or less.
[8] The method for manufacturing a high strength steel sheet described in [7], the method further including, subsequent to the galvanizing treatment, performing an alloying treatment at 450°C or more and 600°C or less. Advantageous Effects of Invention

According to the present invention, a high strength steel sheet that has a TS (tensile strength) of 980 MPa or more and is excellent in terms of formability, that is, specifically, not only ductility but also stretch flangeability and bendability, can be produced. Applying a high strength steel sheet produced by the manufacturing method according to the present invention to, for example, automotive structural components reduces the weights of car bodies and thereby improves fuel economy. Thus, the use of the high strength steel sheet is highly valuable from an industrial viewpoint.

Description of Embodiments

The present invention is specifically described below. Hereinafter, "%" used for describing the contents of elements means "% by mass" unless otherwise specified.
(1) The reasons for which the contents of the constituents of steel are limited to the above ranges in the present invention are described below.

C: 0.030% or More and 0.250% or Less

C is an element necessary for forming tempered martensite and thereby increasing strength. C is also an element effective for enhancing the stability of retained austenite and ductility of steel. If the C content is less than 0.030%, it becomes difficult to achieve the intended area fraction of tempered martensite and the intended strength may fail to be achieved. Moreover, it becomes difficult to achieve a sufficient area fraction of retained austenite and suitable ductility may fail to be achieved.
If C is added to steel in an excessive amount such that the C content exceeds 0.250%, the area fraction of hard tempered martensite is increased to an excessive level. In such a case, during a hole expanding test, the amount of microvoids formed at the grain boundaries of tempered martensite may be increased and the propagation of cracks may proceed disadvantageously. Thus, stretch flangeability may become degraded. Furthermore, the weld zone and the heat-affected zone may become hardened significantly and, consequently, the mechanical properties of the weld zone may become degraded. Thus, spot weldability, arc weldability, and the like may become degraded. From the above-described viewpoints, the C content is limited to 0.030% or more and 0.250% or less. The C content is preferably 0.080% or more and 0.200% or less.

Si: 0.01% or More and 3.00% or Less

Si is effective for achieving suitable ductility since it enhances the strain hardenability of ferrite. If the Si content is less than 0.01%, the advantageous effects of the addition of Si may be poor. Thus, the lower limit is set to 0.01%. Adding Si to steel in an excessive amount such that the Si content exceeds 3.00% may cause steel embrittlement, which degrades ductility and bendability, and formation of red scale, which degrades surface quality. Furthermore, the quality of coating may become degraded. Accordingly, the Si content is limited to 0.01% or more and 3.00% or less. The Si content is preferably 0.20% or more and 2.00% or less and is more preferably 0.20% or more and less than 0.70%.

Mn: 2.50% or More and 8.00% or Less

Mn is an additional element extremely important in the present invention. Mn is an element that stabilizes retained austenite and is effective for achieving suitable ductility. Mn is also an element that increases steel strength by solid solution strengthening. Mn is also effective for forming stable retained austenite at which Mn is concentrated, thereby enabling the intended quotient of the average Mn content (mass%) in retained austenite divided by the average Mn content (mass%) in ferrite to be achieved, and achieving suitable ductility. The above-described actions are confirmed when the Mn content in steel is 2.50% or more. However, if Mn is added to steel in an excessive amount such that the Mn content exceeds 8.00%, a nonuniform microstructure may be formed due to the segregation of Mn and, consequently, stretch flangeability may become degraded. Furthermore, the intended quotient of the average C content (mass%) in retained austenite divided by the average Mn content (mass%) in retained austenite may fail to be achieved. This makes it difficult to achieve both suitable ductility and bendability. Furthermore, ease of conversion treatment and the quality of coating may become degraded. From the above-described viewpoints, the Mn content is limited to 2.50% or more and 8.00% or less. The Mn content is preferably 3.10% or more and 6.00% or less and is more preferably 3.20% or more and 4.20% or less.

P: 0.001% or More and 0.100% or Less

P is an element that causes solid solution strengthening and may be added to steel in accordance with the intended strength. P is also an element that facilitates ferrite transformation and is therefore effective for forming a multiphase microstructure. In order to produce the above-described advantageous effects, it is necessary to limit the P content to 0.001% or more. If the P content exceeds 0.100%, weldability may become degraded. Moreover, in the case where the alloying treatment of zinc coating is performed, the alloying rate may be reduced and the quality of the zinc coating may become degraded. Accordingly, the P content is limited to 0.001% or more and 0.100% or less and is preferably 0.005% or more and 0.050% or less.

S: 0.0001% or More and 0.0200% or Less

S segregates at grain boundaries to cause steel embrittlement during hot working and is present in the form of sulfides, which degrade local deformability. Accordingly, it is necessary to limit the S content to 0.0200% or less. The S content is preferably 0.0100% or less and is more preferably 0.0050% or less. However, due to the limitations of production technology, it is necessary to limit the S content to 0.0001% or more. Accordingly, the S content is limited to 0.0001% or more and 0.0200% or less. The S content is preferably 0.0001% or more and 0.0100% or less and is more preferably 0.0001% or more and 0.0050% or less.

N: 0.0005% or More and 0.0100% or Less

N is an element that degrades the aging resistance of steel. The degradation of aging resistance becomes particularly significant if the N content exceeds 0.0100%. Although it is preferable to minimize the N content, it is necessary to limit the N content to 0.0005% or more due to the limitations of production technology. Accordingly, the N content is limited to 0.0005% or more and 0.0100%. The N content is preferably 0.0010% or more and 0.0070% or less.

Al: 0.001% or More and 2.000% or Less

Al is an element that widens the ferrite-austenite dual phase region and is effective for reducing the dependence of mechanical properties on the annealing temperature, that is, enhancing the stability of mechanical properties. The lower limit is set to 0.001% because the advantageous effects of addition of Al are poor if the Al content is less than 0.001%. Al is also an element that serves as a deoxidizing agent and is effective for the index of cleanliness of steel. Thus, it is preferable to add Al to steel in the deoxidization step. However, adding Al to steel in an excessive amount such that the Al content exceeds 2.000% increases the risk of steel slabs cracking during continuous casting and reduces manufacturability. From the above viewpoints, the Al content is limited to 0.001% or more and 2.000% or less. The Al content is preferably 0.200% or more and 1.200% or less.
In addition to the above-described constituents, the chemical composition may contain at least one element selected from, by mass, Ti: 0.200% or less, Nb: 0.200% or less, V: 0.500% or less, W: 0.500% or less, B: 0.0050% or less, Ni: 1.000% or less, Cr: 1.000% or less, Mo: 1.000% or less, Cu: 1.000% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ta: 0.100% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, and REM (abbreviation for rare earth metals): 0.0050% or less.

Ti: 0.200% or Less

Ti is effective for the precipitation strengthening of steel. Ti increases the strength of ferrite, thereby reduces the difference in hardness between ferrite and the hard second phase (tempered martensite or retained austenite), and consequently enables further suitable stretch flangeability to be achieved. In the case where Ti is added to steel, the Ti content is preferably 0.005% or more and is further preferably 0.010% or more. However, if the Ti content exceeds 0.200%, the area fraction of hard tempered martensite is increased to an excessive level. In such a case, during a hole expanding test, the amount of microvoids formed at the grain boundaries of tempered martensite may be increased and the propagation of cracks may proceed disadvantageously. Thus, stretch flangeability may become degraded. Accordingly, in the case where Ti is added to steel, the Ti content is limited to 0.200% or less. The Ti content is preferably 0.100% or less.

Nb: 0.200% or Less, V: 0.500% or Less, and W: 0.500% or Less

Nb, V, and W are effective for the precipitation strengthening of steel. Furthermore, similarly to the advantageous effects of addition of Ti, Nb, V, and W increase the strength of ferrite, thereby reduce the difference in hardness between ferrite and the hard second phase (tempered martensite or retained austenite), and consequently enable further suitable stretch flangeability to be achieved. In the case where Nb, V, and W are added to steel, the Nb, V, and W contents are each preferably 0.005% or more and are each more preferably 0.010% or more. However, if the Nb content exceeds 0.200% or the V or W content exceeds 0.500%, the area fraction of hard tempered martensite is increased to an excessive level. In such a case, during a hole expanding test, the amount of microvoids formed at the grain boundaries of tempered martensite may be increased and the propagation of cracks may proceed disadvantageously. Thus, stretch flangeability may become degraded. Accordingly, in the case where Nb is added to steel, the Nb content is limited to 0.200% or less. The Nb content is preferably 0.100% or less. In the case where V or W is added to steel, the V or W content is limited to 0.500% or less. The V or W content is preferably 0.300% or less.

B: 0.0050% or Less

B inhibits the formation and growth of ferrite at austenite grain boundaries. B increases the strength of ferrite, thereby reduces the difference in hardness between ferrite and the hard second phase (tempered martensite or retained austenite), and consequently enables further suitable stretch flangeability to be achieved. In the case where B is added to steel, the B content is preferably 0.0003% or more. However, if the B content exceeds 0.0050%, formability may become degraded. Accordingly, in the case where B is added to steel, the B content is limited to 0.0050% or less. The B content is preferably 0.0030% or less.

Ni: 1.000% or Less

Ni is an element that stabilizes retained austenite and is effective for achieving further suitable ductility. Ni is also an element that increases steel strength by solid solution strengthening. In the case where Ni is added to steel, the Ni content is preferably 0.005% or more.
However, if the Ni content exceeds 1.000%, the area fraction of hard tempered martensite is increased to an excessive level. In such a case, during a hole expanding test, the amount of microvoids formed at the grain boundaries of tempered martensite may be increased and the propagation of cracks may proceed disadvantageously. Thus, stretch flangeability may become degraded. Accordingly, in the case where Ni is added to steel, the Ni content is limited to 1.000% or less.

Cr: 1.000% or Less and Mo: 1.000% or Less

Cr and Mo may be added to steel as needed because they enhance the balance between strength and ductility. In the case where Cr and Mo are added to steel, the Cr and Mo contents are preferably Cr: 0.005% or more and Mo: 0.005% or more. However, if Cr and Mo are added to steel in excessive amounts such that the Cr content exceeds 1.000% and the Mo content exceeds 1.000%, the area fraction of hard tempered martensite is increased to an excessive level. In such a case, during a hole expanding test, the amount of microvoids formed at the grain boundaries of tempered martensite may be increased and the propagation of cracks may proceed disadvantageously. Thus, stretch flangeability may become degraded. Accordingly, in the case where the above elements are added to steel, the contents of the elements are limited to Cr: 1.000% or less and Mo: 1.000% or less.

Cu: 1.000% or Less

Cu may be added to steel as needed because it is an element effective for strengthening of steel. If Cu is added to steel, the Cu content is preferably 0.005% or more. However, if Cu is added to steel in an excessive amount such that the Cu content exceeds 1.000%, the area fraction of hard tempered martensite is increased to an excessive level. In such a case, during a hole expanding test, the amount of microvoids formed at the grain boundaries of tempered martensite may be increased and the propagation of cracks may proceed disadvantageously. Thus, stretch flangeability may become degraded. Accordingly, in the case where Cu is added to steel, the Cu content is limited to 1.000% or less.

Sn: 0.200% or Less and Sb: 0.200% or Less

Sn and Sb are added to steel as needed in order to inhibit a region of the surface layer of the steel sheet which has a thickness of several tens of micrometers from being decarburized as a result of nitriding or oxidation of the surface of the steel sheet. Sn and Sb are effective for inhibiting the above nitriding or oxidation, preventing a reduction in the area fraction of tempered martensite in the surface of the steel sheet, and achieving certain strength and the stability of mechanical properties. In the case where Sn and Sb are added to steel, the Sn and Sb contents are each preferably 0.002% or more. However, if any of the above elements is added to steel in an excessive amount such that the content of the element exceeds 0.200%, toughness may become degraded. Accordingly, in the case where Sn and Sb are added to steel, the Sn and Sb contents are each limited to 0.200% or less.

Ta: 0.100% or Less

Similarly to Ti and Nb, Ta increases strength by forming an alloy carbide and an alloy carbonitride. Furthermore, it is considered that Ta partially dissolves in Nb carbide or carbonitride to form a composite precipitate, such as (Nb,Ta)(C,N), thereby markedly reduces the coarsening of the precipitates, and stabilizes the increase in strength by precipitation strengthening. Therefore, Ta is preferably added to steel. In the case where Ta is added to steel, the Ta content is preferably 0.001% or more. Even if Ta is added to steel in an excessive amount, the above precipitation stabilization effect may become saturated and the costs of the alloying elements are also increased. Accordingly, in the case where Ta is added to steel, the Ta content is limited to 0.100% or less.

Ca: 0.0050% or Less, Mg: 0.0050% or Less, Zr: 0.0050% or Less, and REM: 0.0050% or Less

Ca, Mg, Zr, and REMs are elements effective for increasing the sphericity of sulfides and reducing the adverse effects of the sulfides on stretch flangeability. In the case where the above elements are added to steel, the contents of the elements are each preferably 0.0005% or more. However, if any of the above elements is added to steel in an excessive amount such that the content of the element exceeds 0.0050%, the amount of inclusions and the like may be increased and, consequently, surface and internal defects may occur. Accordingly, if Ca, Mg, Zr, and REMs are added to steel, the Ca, Mg, Zr, and REM contents are each limited to 0.0050% or less.
The constituent other than the above-described constituents, that is, the balance, includes Fe and incidental impurities.
(2) The steel microstructure is described below.

Area Fraction of Ferrite: 30% or More and 80% or Less

It is necessary to limit the area fraction of ferrite to 30% or more in order to achieve sufficient ductility. It is also necessary to limit the area fraction of soft ferrite to 80% or less in order to achieve a tensile strength of 980 MPa or more. Note that the term "ferrite" used herein refers to polygonal ferrite, granular ferrite, and acicular ferrite, that is, ferrite materials that are relatively soft and rich in ductility. The area fraction of ferrite is preferably 40% or more and 75% or less.

Area Fraction of Tempered Martensite: 3.0% or More and 35% or Less

The area fraction of tempered martensite needs to be 3.0% or more in order to achieve a high local elongation, suitable stretch flangeability, suitable bendability, and a high yield ratio. It is necessary to limit the area fraction of tempered martensite to 3.0% or more in order to achieve a high local elongation, suitable stretch flangeability, and suitable bendability. It is also necessary to limit the area fraction of tempered martensite to 35% or less because ductility and bendability may become degraded if the fraction of tempered martensite is excessively high. The area fraction of tempered martensite is preferably 5.0% or more and 20% or less.
The area fractions of ferrite and tempered martensite can be determined by grinding a thickness-wise cross section (L-section) of the steel sheet which is parallel to the rolling direction, etching the L-section in 3-vol% nital, then observing the 1/4-thickness position (i.e., the position 1/4 of the thickness of the steel sheet below the surface in the depth direction) of the L-section with a SEM (scanning electron microscope) at a 2000-fold magnification for 10 fields of view, calculating the area fractions of the microstructure components (ferrite and tempered martensite) with "Image-Pro" produced by Media Cybernetics, Inc. for each of the 10 fields of view on the basis of the resulting microstructure images, and taking the averages thereof. In the above microstructure images, ferrite appears as a gray microstructure component (base microstructure) and tempered martensite appears as a microstructure component that includes a white (As-quenched) martensite portion and a gray inner structure formed inside the martensite portion.

Area Fraction of Retained Austenite: 8% or More

It is necessary to limit the area fraction of retained austenite to 8% or more in order to achieve sufficient ductility. The area fraction of retained austenite is preferably 12% or more and 25% or less.
The area fraction of retained austenite was determined by grinding the steel sheet to a position 0.1 mm below the 1/4-thickness position, further grinding the steel sheet 0.1 mm by chemical polishing, then measuring the integrated intensity ratios of diffraction peaks on the {200}, {220}, and {311} planes of fcc iron and the {200}, {211}, and {220} planes of bcc iron with an X-ray diffraction apparatus using CoKα radiation, and taking the averages of the nine integrated intensity ratios.

Quotient of Area Fraction of Retained Austenite Grains Having Aspect Ratio of 2.0 or More and Minor Axis Length of 1 µm or Less Divided by Total Area Fraction of Retained Austenite Is 0.3 or More

Limiting the quotient of the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less divided by the total area fraction of the retained austenite to 0.3 or more is an important condition constituting the present invention. The retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less reduce the likelihood of voids being formed during punching performed prior to the flange forming step and thereby enhance stretch flangeability. For achieving suitable stretch flangeability, the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less needs to be high, in addition to satisfying the area fraction of retained austenite which is high enough to achieve high ductility. The above index is preferably 0.5 or more. The upper limit for the above aspect ratio is preferably 15.0 or less. The lower limit for the above minor axis length is preferably 0.05 µm or more, which is the detection limit in EBSD.
Martensite and retained austenite were identified using a phase map obtained by EBSD (electron backscattered diffraction). The aspect ratio of a retained austenite grain was calculated by drawing an ellipse circumscribing the retained austenite grain with "Photoshop elements 13" and dividing the length of the major axis of the ellipse by the length of the minor axis of the ellipse.

Quotient of Average Mn Content (Mass%) in Retained Austenite Divided by Average Mn Content (Mass%) in Ferrite: 1.5 or More

Limiting the quotient of the average Mn content (mass%) in retained austenite divided by the average Mn content (mass%) in ferrite to 1.5 or more is an extremely important condition constituting the present invention. For achieving suitable ductility, the area fraction of stable retained austenite at which Mn is concentrated needs to be high. The above index is preferably 2.0 or more. Although the upper limit for the above quotient is not limited because the higher the average Mn content in retained austenite, the higher the ductility, the above index is preferably 10.0 or less because the increase in ductility becomes saturated if the above index exceeds 10.0.

Product of Quotient of Average Mn Content (Mass%) in Retained Austenite Divided by Average Mn Content (Mass%) in Ferrite and Average Aspect Ratio of Retained Austenite Grains Is 3.0 or More

Limiting the product of the quotient of the average Mn content (mass%) in retained austenite divided by the average Mn content (mass%) in ferrite and the average aspect ratio of retained austenite grains to 3.0 or more is an extremely important condition. For achieving suitable ductility, the area fraction of stable retained austenite grains having a high aspect ratio at which Mn is concentrated needs to be high. The above index is preferably 4.0 or more. The upper limit for the above index is suitably 20.0 or less.

Quotient of C Content (Mass%) in Retained Austenite Divided by Average C Content (Mass%) in Ferrite: 3.0 or More

Limiting the quotient of the C content (mass%) in retained austenite divided by the average C content (mass%) in ferrite to 3.0 or more is an extremely important condition constituting the present invention. When the C content in retained austenite is increased, retained austenite remains even in high strain region during bending deformation and, consequently, bendability is enhanced. Therefore, for achieving suitable ductility and bendability, the area fraction of stable retained austenite at which Mn is concentrated needs to be high. The above index is preferably 5.0 or more. The upper limit for the above index is preferably 10.0 or less.

Quotient of Average C Content (Mass%) in Retained Austenite Divided by Average Mn Content (Mass%) in Retained Austenite Is 0.05 or More

Limiting the quotient of the C content (mass%) in retained austenite divided by the Mn content (mass%) in retained austenite to 0.05 or more is an extremely important condition constituting the present invention. For achieving suitable bendability, the area fraction of stable retained austenite at which not only Mn but also C are concentrated needs to be high. When not only Mn but also C are concentrated at retained austenite, stable retained austenite can be present in high strain region and, consequently, compatibility of suitable ductility and bendability can be achieved. The above index is preferably 0.06 or more. The upper limit for the above index is preferably 0.10 or less.
The C and Mn contents in retained austenite and ferrite are determined with 3-dimensional atom probe (3DAP) using a sample taken from the 1/4-thickness position. First, a portion of the steel sheet which includes retained austenite and ferrite is sampled and then formed into an acicular sample using focused ion beam. A voltage is applied to the acicular sample with the 3DAP, and C and Mn ions released upon the voltage application are analyzed. The Mn content can be determined in terms of atom% by dividing the numbers of C and Mn atoms measured by the total number of all the atoms measured, for each of retained austenite and ferrite. The above measurement is performed for each of 30 retained austenite grains and 30 ferrite grains which are randomly selected in the field of view, and the averages of the C and Mn contents determined by the quantitative analysis are calculated. The C and Mn contents (mass%) in retained austenite and ferrite can be obtained by converting the C and Mn contents (atom%) to a mass basis.

Content of Diffusible Hydrogen in Steel Is 0.3 ppm by Mass or Less

Limiting the content of diffusible hydrogen in steel to 0.3 ppm by mass or less is an important condition constituting the present invention. For achieving a high local elongation and suitable stretch flangeability, it is necessary to limit the content of diffusible hydrogen in steel to 0.3 ppm by mass or less. The content of diffusible hydrogen in steel is preferably 0.2 ppm by mass or less. A test specimen having a length of 30 mm and a width of 5 mm was taken from an annealed steel sheet. After the coated layer had been removed by grinding, the content of diffusible hydrogen in steel and the peak of release of diffusible hydrogen were measured. The release peak was measured by thermal desorption spectrometry (TDS) at a heating rate of 200 °C/hr. The hydrogen detected at 300°C or less was considered diffusible hydrogen.
The advantageous effects of the present invention are not impaired even if the steel microstructure according to the present invention contains quenched martensite, bainite, pearlite, or carbides, such as cementite, in addition to ferrite, tempered martensite, and retained austenite, when the area fraction of the quenched martensite, bainite, pearlite, or carbides is 10% or less.
The steel sheet may include a galvanized layer disposed on the surface. The galvanized layer may be a galvannealed layer formed by performing an alloying treatment of the galvanized layer.
(3) Manufacturing conditions are described below.

Temperature at Which Steel Slab Is Heated

The temperature at which the slab is heated is preferably, but not limited to, 1100°C or more and 1300°C or less. Since the precipitates present when the steel slab is heated are present in the form of coarse precipitates in the final steel sheet and do not affect strength, the Ti- and Nb-based precipitates formed during casting can be again dissolved. It is also preferable to limit the temperature at which the steel slab is heated to 1100°C or more in order to scale off air bubbles, segregation, and the like present in the slab surface layer, further reduce cracks and irregularities present in the surface of the steel sheet, and thereby further flatten the surface of the steel sheet. The temperature at which the steel slab is heated is preferably 1300°C or less in order to reduce scale loss caused with an increase in oxidation. The above slab heating temperature is more preferably 1150°C or more and 1250°C or less.
The steel slab is preferably manufactured by a continuous casting process in order to prevent macrosegregation. Alternatively, the steel slab may be manufactured by ingot casting, thin-slab casting, or the like. In addition to the conventional method in which, subsequent to the production of the steel slab, the steel slab is cooled to room temperature and then reheated, energy-saving processes, such as hot charge rolling and hot direct rolling, in which the steel slab is not cooled to room temperature but charged into a heating furnace while the temperature of the steel slab is high or the steel slab is rolled immediately after heat insulation has been done for a short period of time, may be used without problems. The slab is formed into a sheet bar by rough rolling under common conditions. In the case where the heating temperature is relatively low, the sheet bar is preferably heated with a bar heater or the like prior to finish rolling in order to prevent the occurrence of troubles during hot rolling.

Finish Rolling Delivery Temperature in Hot Rolling: 750°C or More and 1000°C or Less

The heated steel slab is hot rolled to form a hot rolled steel sheet by rough rolling and finish rolling. If the above finish rolling delivery temperature exceeds 1000°C, the amount of oxides (scales) formed rapidly increases, the roughness of the interface between the base iron and the oxides is increased accordingly, and surface quality may become degraded subsequent to pickling and cold rolling. Furthermore, if hot rolling scales partially remain unremoved subsequent to pickling, they adversely affect ductility and stretch flangeability. Moreover, grain size may be increased to an excessive level and the roughness of the surface of a pressed article may be increased during working. If the above finish rolling delivery temperature is less than 750°C, the rolling force is increased, the rolling load is increased accordingly, and the rolling reduction ratio performed while austenite is in a non-recrystallization state is increased. In such a case, an abnormal texture develops. Consequently, the in-plane anisotropy of the final product is increased significantly. This degrades the uniformity of material quality (stability of mechanical properties). In addition, the aspect ratio of retained austenite grains is reduced and, consequently, ductility and stretch flangeability may become degraded. Accordingly, it is necessary to limit the finish rolling delivery temperature in hot rolling to 750°C or more and 1000°C or less. The finish rolling delivery temperature is preferably 800°C or more and 950°C or less.

Temperature at Which Coiling Is Performed Subsequent to Hot Rolling: 300°C or More and 750°C or Less

If the temperature at which coiling is performed subsequent to hot rolling exceeds 750°C, the grain size of ferrite included in the microstructure of the hot rolled steel sheet is increased, the aspect ratio of retained austenite grains included in the final annealed steel sheet is reduced, it becomes difficult to achieve the intended quotient of the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less divided by the total area fraction of the retained austenite, and stretch flangeability becomes degraded consequently. If the temperature at which coiling is performed subsequent to hot rolling is less than 300°C, the strength of the hot rolled steel sheet is increased. In such a case, the rolling load required for cold rolling may be increased and defects of the shape of the steel sheet may occur. This reduces productivity. Accordingly, it is necessary to limit the temperature at which coiling is performed subsequent to hot rolling to 300°C or more and 750°C or less. The above coiling temperature is preferably 400°C or more and 650°C or less.
The finish rolling may be performed in a continuous manner by joining the rough rolled steel sheets to one another during hot rolling. The rough rolled steel sheets may be coiled temporarily. In order to reduce the rolling force required for hot rolling, a part or the entirety of the finish rolling may be performed using a lubricant. It is also preferable to perform lubrication rolling in order to increase the uniformity in the shape of the steel sheet and the uniformity of material quality. In the case where lubrication rolling is performed, the coefficient of friction at which the lubrication rolling is performed is preferably 0.10 or more and 0.25 or less.
The hot rolled steel sheet produced in the above-described manner may be pickled optionally. It is preferable to perform pickling because it enables the removal of oxides present on the surface of the steel sheet and further improves ease of conversion treatment and the quality of coating. The pickling may be performed in a single stage or multiple stages.

Holding at Temperature Equal to or Lower Than Ac₁ Transformation Temperature for More Than 1800 s

It is preferable to perform holding at a temperature equal to or lower than the Ac₁ transformation temperature for more than 1800 s because it softens the steel sheet that is to be cold rolled in the subsequent step. If holding is performed at a temperature higher than the Ac₁ transformation temperature, Mn concentrates at austenite, hard martensite and retained austenite are formed subsequent to cooling, and the steel sheet may fail to be softened. In such a case, in the subsequent annealing step, retained austenite may be formed at the grain boundaries and the amount of retained austenite grains having a low aspect ratio may be increased. This makes it difficult to achieve the intended quotient of the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less divided by the total area fraction of the retained austenite and may consequently degrade stretch flangeability. Even at a temperature equal to or lower than the Ac₁ transformation temperature, if the holding time is 1800 s or less, it becomes difficult to eliminate the strain that remains subsequent to hot rolling and the steel sheet may fail to be softened.
The heat treatment may be performed by any annealing method, such as continuous annealing or batch annealing. Subsequent to the above heat treatment, cooling is performed to room temperature. The method with which the cooling is performed and the rate at which the cooling is performed are not limited. Any cooling method, such as furnace cooling or natural cooling in batch annealing, or gas jet cooling, mist cooling, or water cooling in continuous annealing, may be used. In the case where a pickling treatment is performed, a conventional method may be used.

Cold Rolling

The resulting steel sheet is cold rolled. The cold rolling reduction ratio is preferably, but not limited to, 15% or more and 80% or less. Performing cold rolling at the above reduction ratio enables formation of an sufficiently recrystallized, intended microstructure and further improves ductility.

Holding at Temperature Equal to or Higher Than Ac₃ Transformation Temperature - 50°C for 20 s or More and 1800 s or Less

If holding is performed at a temperature less than the Ac₃ transformation temperature - 50°C, Mn concentrates at austenite, martensitic transformation does not occur during cooling, and cores of retained austenite grains having a high aspect ratio may fail to be formed consequently. In such a case, in the subsequent annealing step, retained austenite may be formed at the grain boundaries disadvantageously. This increases the amount of retained austenite grains having a low aspect ratio and makes it impossible to form the intended microstructure. Even at a temperature equal to or higher than the Ac₃ transformation temperature - 50°C, if the holding time is less than 20 s, recrystallization does not occur to a sufficient degree and the intended microstructure fails to be formed. Consequently, ductility may become degraded. If the above holding is performed for more than 1800 s, the amount of hard tempered martensite is increased subsequent to the final annealing treatment and ductility and bendability may become degraded consequently.

Cooling to Cooling Stop Temperature Equal to or Lower Than Martensitic Transformation Start Temperature

If the cooling stop temperature is higher than the martensitic transformation start temperature, when the amount of martensite that is to be transformed is small, all of the untransformed austenite may transform to martensite during the final cooling, and cores of retained austenite grains having a high aspect ratio fail to be formed. In such a case, in the subsequent annealing step, retained austenite may be formed at the grain boundaries disadvantageously. This increases the amount of retained austenite grains having a low aspect ratio and makes it impossible to form the intended microstructure. The cooling stop temperature is preferably equal to or higher than the martensitic transformation start temperature - 250°C and equal to or lower than the martensitic transformation start temperature - 50°C.

Reheating to Reheating Temperature of 120°C or More and 450°C or Less and Holding at Reheating Temperature for 2 s or More and 600 s or Less

If the reheating temperature is less than 120°C, C is not concentrated at retained austenite formed in the subsequent annealing step and the intended microstructure fails to be formed consequently. If the reheating temperature exceeds 450°C, the decomposition of cores of retained austenite grains having a high aspect ratio occurs and the amount of retained austenite grains having a low aspect ratio is increased accordingly. This makes it impossible to form the intended microstructure. If the holding time is less than 2 s, similarly, the cores of retained austenite grains having a high aspect ratio cannot be formed and the intended microstructure fails to be formed. If the holding time exceeds 600 s, the decomposition of cores of retained austenite grains having a high aspect ratio occurs and the amount of retained austenite grains having a low aspect ratio is increased accordingly. This makes it impossible to form the intended microstructure. Cooling is performed to room temperature subsequent to the reheating and holding. The cooling method is not limited. For example, it is preferable to perform natural cooling, gas cooling, or mist cooling to 200°C or less and then perform water cooling. Pickling may be performed as needed. In the case where pickling is performed, common picking methods may be used.

Holding at Temperature Equal to or Higher Than Ac₁ Transformation Temperature and Equal to or Lower Than Ac₁ Transformation Temperature + 150°C for 20 s or More and 600 s or Less

Performing holding at a temperature equal to or higher than the Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less is an extremely important condition constituting the present invention. If the holding is performed at a temperature lower than the Ac₁ transformation temperature for less than 20 s, the carbides formed during heating remain undissolved, it becomes difficult to form sufficient area fractions of tempered martensite and retained austenite, and strength is reduced consequently. If the above holding temperature exceeds the Ac₁ transformation temperature + 150°C, the area fraction of martensite is increased subsequent to the annealing treatment, the amount of tempered martensite present subsequent to the final annealing treatment is increased to an excessive level, and the intended microstructure fails to be formed consequently. Furthermore, Mn and C fail to concentrate at austenite at sufficient levels and, consequently, ductility and bendability become degraded. In addition, a sufficient area fraction of retained austenite fails to be formed and ductility becomes degraded consequently. The above holding temperature is preferably equal to or lower than the Ac₁ transformation temperature + 100°C. If the holding time exceeds 600 s, the area fraction of martensite is increased subsequent to the annealing treatment, the amount of tempered martensite present subsequent to the final annealing treatment is increased to an excessive level, and the intended microstructure fails to be formed consequently. Furthermore, Mn and C fail to concentrate at austenite at sufficient levels and, consequently, ductility and bendability become degraded. In addition, the area fraction of retained austenite which is enough to achieve the intended ductility fails to be formed.

Cooling to 100°C or Less

Performing cooling to 100°C or less enables the amount of martensite which is large enough to achieve the certain tensile strength to be formed. The cooling stop temperature is preferably about 20°C or more and about 50°C or less in consideration of the manufacturing costs.

Galvanizing Treatment

In the case where a galvanizing treatment is performed, the galvanizing treatment is performed before the holding is performed at the temperature equal to or higher than the Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less and then the cooling is performed to 100°C or less. Since this cooling treatment is performed in order to increase coatability by controlling the temperature of the steel sheet that is to be dipped into the plating bath, the cooling is performed to a temperature equal to or higher than the temperature of the galvanizing bath and equal to or lower than about 600°C or less. The cooling method is not limited. For example, natural cooling, furnace cooling, and gas jet cooling may be used. In the case where a hot-dip galvanizing treatment is performed, the steel sheet that has been subjected to the above annealing treatment is dipped into a galvanizing bath having a temperature of 440°C or more and 500°C or less in order to perform a hot-dip galvanizing treatment. Then, the weight of coating is adjusted by gas wiping or the like. For hot-dip galvanizing, a galvanizing bath having an Al content of 0.08% or more and 0.30% or less is preferably used. Instead of a hot-dip galvanizing treatment, an electrogalvanizing treatment or the like may be used.
In the case where an alloying treatment of zinc coating is performed, subsequent to the galvanizing treatment, an alloying treatment of zinc coating is performed at 450°C or more and 600°C or less. If the alloying treatment is performed at more than 600°C, untransformed austenite may transform into pearlite, and the intended area fraction of retained austenite may fail to be achieved. Consequently, ductility may become degraded. Accordingly, in the case where an alloying treatment of zinc coating is performed, it is preferable to perform an alloying treatment of zinc coating at 450°C or more and 600°C or less.
Although the other conditions for the manufacturing method are not limited, it is preferable to perform the above annealing treatment using a continuous annealing facility from the viewpoint of productivity. It is also preferable to perform the set of treatments, such as annealing, hot-dip galvanizing, and alloying of zinc coating, using CGL (continuous galvanizing line), which is a hot-dip galvanizing line.

Holding at More Than 100°C and 400°C or Less for 10 s or More

Performing holding at more than 100°C and 400°C or less for 10 s or more as a final heat treatment (final annealing treatment) is important in the present invention. If holding is performed at 100°C or less or for less than 10 s, a sufficient area fraction of tempered martensite fails to be formed. Furthermore, diffusible hydrogen included in the steel is not released from the steel sheet. This results in degradation of stretch flangeability and bendability. If holding is performed at more than 400°C, a sufficient area fraction of retained austenite fails to be formed due to the decomposition of retained austenite. Consequently, steel ductility becomes degraded.
The above "high strength steel sheet" and the "high strength galvanized steel sheet" that includes the high strength steel sheet and a galvanized layer disposed on the surface may be subjected to skin pass rolling in order to, for example, perform shape correction and the adjustment of surface roughness. The rolling reduction ratio in the skin pass rolling is preferably 0.1% or more and 2.0% or less. If the above rolling reduction ratio is less than 0.1%, the advantageous effects are small and it is difficult to control the reduction ratio. Thus, this is considered the lower limit for the suitable rolling reduction ratio. If the rolling reduction ratio exceeds 2.0%, productivity may become degraded significantly. Thus, this is considered the upper limit for the suitable rolling reduction ratio. The skin pass rolling may be performed either online or offline. Skin pass rolling at an intended rolling reduction ratio may be performed in a single stage or multiple stages. Furthermore, various coating treatments, such as resin coating and fat coating, may be used.

EXAMPLES

Molten steels having the chemical compositions described in Table 1 with the balance being Fe and incidental impurities were prepared with a converter and formed into slabs by continuous casting process. The slabs were reheated to 1250°C and then formed into high strength cold rolled steel sheets (CR) having a thickness of 1.0 mm or more and 1.8 mm or less under the conditions described in Tables 2 and 3. Furthermore, a galvanizing treatment was performed to produce hot-dip galvanized steel sheets (GI). An alloying treatment of the hot-dip galvanized steel sheets was performed to produce hot-dip galvannealed steel sheets (GA). The hot-dip galvanizing bath used for producing the hot-dip galvanized steel sheets (GI) was a zinc bath containing Al: 0.19% by mass. The hot-dip galvanizing bath used for producing the hot-dip galvannealed steel sheets (GA) was a zinc bath containing Al: 0.14% by mass. The bath temperature was 465°C. The coating weight was 45 g/m² per side (both sides were coated). In the production of the GA steel sheets, the Fe concentration in the coated layer was adjusted to 9% by mass or more and 12% by mass or less. The cross-sectional steel microstructure, tensile properties, and stretch flangeability of each of the above steel sheets were determined. Tables 4, 5, and 6 list the results.

The martensitic transformation start temperature, that is, the Ms temperature, and the Ac₁ and Ac₃ transformation temperatures were calculated using the following formulae.

\begin{array}{l} Martensitic transformation start temperature: Ms \\ temperature (° C) \\ = 550 - 350 \times (%C) - 40 \times (%Mn) - 10 (%Cu) - 17 \times \\ (%Ni) - 20 \times (%Cr) - 10 \times (%Mo) - 35 \times (%V) - 5 \times (%W) + 30 \times \\ (%Al) \end{array}

\begin{array}{l} {Ac}_{1} Transformation temperature (° C) \\ = 751 - 16 \times (%C) + 11 \times (%Si) - 28 \times (%Mn) - 5.5 \times \\ (%Cu) - 16 \times (%Ni) + 13 \times (%Cr) + 3.4 \times (%Mo) \end{array}

\begin{array}{l} {Ac}_{3} Transformation temperature (° C) \\ = 910 - 203 \sqrt (%C) + 45 \times (%Si) - 30 \times (%Mn) - 20 \times (%Cu) \\ - 15 \times (%Ni) + 11 \times (%Cr) + 32 \times (%Mo) + 104 \times (%V) + 400 \times \\ (%Ti) + 200 \times (%Al) \end{array}

where (%C), (%Si), (%Mn), (%Ni), (%Cu), (%Cr), (%Mo), (%V), (%Ti), (%V), (%W), and (%Al) each represent the content (mass%) of the element.

[Table 2]

No.	Steel type	Finish rolling delivery temperature	Coiling temperature	Heat treatment of hot rolled steel sheet		Cold rolling reduction ratio	Annealing treatment of cold rolled steel sheet					Annealing treatment of cold rolled steel sheet		Alloying temperature	Final annealing treatment		Type'
		Finish rolling delivery temperature	Coiling temperature	Heat treatment temperature	Heat treatment time	Cold rolling reduction ratio	Heat treatment temperature	Heat treatment time	Cooling stop temperature	Reheating temperature	Reheating temperature holding time	Heat treatment temperature	Heat treatment time	Alloying temperature	Heat treatment temperature	Heat treatmen t time
		(°C)	(°C)	(°C)	(s)	(%)	(°C)	(s)	(°C)	(°C)	(s)	(°C)	(s)	(°C)	(°C)	(s)
1	A	890	530	550	21600	66.7	800	135	130	270	180	695	140		120	24400	CR
2	A	870	520	525	19800	58.8	850	160	170	300	350	700	170		180	5400	GI
3	A	860	500	590	32400	50.0	820	140	150	275	120	690	160	520	150	32400	GA
4	A	890	470	540	23400	64.7	815	120	120	300	250	710	310	530	200	3600	GA
5	A	900	530	510	18000	62.5	600	150	200	350	170	670	130		180	64800	GI
6	A	930	510	610	21600	58.8	910	15	215	280	140	700	130	550	310	10800	GA
7	A	860	540	615	14400	50.0	780	2400	80	150	220	710	125	500	320	14400	GA
8	A	800	550	570	32400	52.9	740	200	410	440	80	700	150		220	60000	GI
9	A	820	540			56.3	790	220	125	480	215	730	90	460	130	24400	GA
10	A	860	440	600	18000	64.7	810	120	70	110	320	660	530	480	330	7200	GA
11	A	900	530	510	4400	61.1	795	50	120	190	1000	690	450		210	9000	GI
12	A	885	370	535	12600	66.7	820	360	200	270	1	700	180		140	21600	CR
13	A	940	490			46.2	845	250	180	275	400	690	120	545	170	25200	GA
14	B	900	490	570	21600	54.8	830	800	315	420	540	700	160		200	28800	GI
15	C	900	520	550	18000	52.9	840	150	250	300	180	710	150	520	220	600	GA
16	A	690	530	560	14400	47.1	880	180	140	215	210	730	240		210	14400	GI
17	A	850	840	590	32400	47.8	860	240	120	240	120	665	180	550	150	27000	GA
18	A	870	610	740	25200	58.8	870	180	110	200	50	710	150		370	60	GI
19	A	860	550	430	7200	58.8	790	240	125	195	370	510	300	510	300	23400	GA
20	A	860	540	530	14400	68.4	780	300	175	330	530	880	480	485	160	25200	GA
21	A	900	550	530	10800	72.2	770	150	150	170	180	685	5	525	310	18000	GA
22	A	860	580	540	19800	64.7	740	180	130	250	260	715	2200	530	390	14400	GA
23	A	890	585	560	3600	61.1	800	90	150	225	120	700	130	490	30	57600	GA
24	A	840	590			56.3	820	120	210	310	150	720	140	550	520	18000	GA
25	A	925	540			64.7	815	180	200	300	50	680	100	500	250	5	GA
26	D	895	560	570	28800	58.8	880	360	140	250	180	705	110		370	2400	CR
27	E	950	550	570	18000	58.8	800	1200	180	310	240	690	480		355	10800	GI
28	F	865	600	530	14400	42.9	820	150	175	295	180	770	65		255	36000	CR
29	G	880	620	490	27000	57.1	830	140	100	250	130	710	260	530	260	12000	GA
30	H	885	500	535	36000	53.3	830	120	150	230	270	690	90	510	320	12000	GA
31	I	900	550	530	23400	50.0	765	150	140	340	540	710	300	525	350	72000	GA
32	J	880	500	510	25200	64.7	770	180	130	300	30	690	210		300	4800	CR
33	K	870	410	510	32400	48.6	790	90	65	200	220	660	120		120	96000	GI

Underlined part: Outside of the scope of the present invention
* CR: Cold rolled steel sheet (without coating), GI: Hot-dip galvanized steel sheet (without alloying of zinc coating), GA: Hot-dip galvannealed steel sheet

[Table 3]

No.	Steel type	Finish rolling delivery temperature	Coiling temperature	Heat treatment of hot rolled steel sheet		Cold rolling reduction ratio	Annealing treatment of cold rolled steel sheet					Annealing treatment of cold rolled steel sheet		Alloying temperatur e	Final annealing treatment		Type'
		Finish rolling delivery temperature	Coiling temperature	Heat treatment temperature	Heat treatment time	Cold rolling reduction ratio	Heat treatment temperature	Heat treatment time	Cooling stop temperature	Reheating temperature	Reheating temperature holding time	Heat treatment temperature	Heat treatment time	Alloying temperatur e	Heat treatment temperatur e	Heat treatme nt time
		(°C)	(°C)	(°C)	(s)	(%)	(°C)	(s)	(°C)	(°C)	(s)	(°C)	(s)	(°C)	(°C)	(s)
34	L	960	570	560	19800	46.2	790	100	180	240	150	720	580		170	188000	GI
35	M	890	620	580	39600	62.5	710	150	130	280	150	695	180	540	380	240	GA
36	N	880	600	520	28800	62.5	790	120	130	290	180	705	120		290	24000	CR
37	O	880	530	520	7200	58.8	920	130	250	355	570	770	240	560	230	9000	GA
38	P	860	490			52.0	980	140	150	200	400	790	360	510	240	9600	GA
39	Q	860	560	590	9000	56.3	990	85	210	325	80	780	140		270	14400	CR
40	R	890	525			58.8	910	150	200	350	100	680	180	530	200	10800	GA
41	S	890	640	540	10800	62.5	880	600	130	200	200	720	110	480	270	72000	GA
42	S	870	610	550	10800	62.5	890	720	250	300	180	700	100		340	1200	CR
43	T	860	480	480	10800	64.7	800	60	50	180	190	660	300		120	60000	GI
44	U	895	550	550	36000	57.1	890	110	220	300	140	725	150		210	5400	GI
45	V	860	600	550	32400	50.0	860	330	200	400	240	710	180		150	54000	CR
46	W	900	500			56.3	900	125	220	320	100	695	300	530	200	96000	GA
47	X	910	560	510	36000	46.2	850	150	180	320	225	730	180	520	260	4200	GA
48	Y	870	550	580	10800	52.9	840	140	55	330	170	695	600		160	12000	CR
49	Z	895	330			47.1	815	300	180	300	510	690	240	510	290	72000	GA
50	AA	880	600	540	32400	55.6	830	900	140	405	260	700	90	540	200	4800	GA
51	AB	900	530	540	14400	56.3	890	140	150	180	240	720	180		270	96000	GI
52	AC	880	740	520	19800	58.8	790	60	120	250	160	710	250		120	188000	GI
53	AD	910	600	520	21600	53.3	890	220	150	300	90	720	180	525	120	24000	GA
54	AE	880	500	530	27000	64.7	810	120	140	325	210	700	150		190	10800	GI
55	AF	900	510	560	9000	62.5	840	180	210	320	540	710	60	550	210	24000	GA
56	AG	910	580	520	28800	56.3	830	150	200	320	210	730	45	500	150	28800	GA
57	AH	845	570			61.5	810	160	150	200	180	740	250	510	230	18000	GA
58	AI	900	560	530	32400	56.3	900	120	90	180	80	725	180		310	2400	GI
59	AJ	910	560	540	14400	56.3	910	80	100	200	125	690	150	520	190	27000	GA
60	AK	860	550	540	14400	56.3	800	90	180	210	150	715	240		180	36000	CR
61	AL	870	520	510	7200	64.7	825	150	170	370	150	700	120	480	255	188000	GA
62	AM	830	500			46.7	840	180	100	300	240	700	90		260	25200	CR
63	AN	850	500	560	21600	50.0	830	140	95	300	60	685	210	550	370	30	GA
64	AO	890	510	515	12600	57.1	825	240	120	380	320	690	220	540	140	9600	GA

The tensile test was conducted in accordance with JIS Z 2241 (2011) using a JIS No. 5 test specimen prepared by taking a sample from each of the steel sheets such that the tensile direction was perpendicular to the rolling direction of the steel sheet. In the tensile test, TS (tensile strength) and EL (total elongation) were measured. As for mechanical properties, an evaluation of "Good" was given in the following cases.

When TS was 980 MPa or more and less than 1080 MPa, EL ≥ 20%
When TS was 1080 MPa or more and less than 1180 MPa, EL ≥ 16%

Stretch flangeability was measured in accordance with JIS Z 2256 (2010). Specifically, each of the steel sheets was cut into a piece having a size of 100 mm × 100 mm. A hole having a diameter of 10 mm was formed in the piece at a clearance of 12% ± 1%. Subsequently, while the piece was held by a die having an inside diameter of 75 mm at a blank holding force of 9 ton, a conical punch having a 60 degree apex angle was inserted into the hole. Then, the diameter of the hole was measured at the critical point at which cracking occurred. A critical hole expanding ratio λ(%) was calculated using the following formula. The evaluation of stretch flangeability was made on the basis of the critical hole expanding ratio. $Critical hole expanding ratio λ (%) = \{(D_{f} - D_{0}) / D_{0}\} \times 100$
where D_f represents the diameter (mm) of the hole at which cracking occurred, and D₀ represents the initial diameter (mm) of the hole. In the present invention, an evaluation of "Good" was given in the following cases depending on TS.

When TS was 980 MPa or more and less than 1080 MPa, λ ≥ 20%
When TS was 1080 MPa or more and less than 1180 MPa, λ ≥ 15%

The bend test was conducted in accordance with the V block method described in JIS Z 2248 (1996) using a bend test specimen having a width of 30 mm and a length of 100 mm which was taken from each of the annealed steel sheets such that the rolling direction was parallel to the bending axis (bending direction). The test was conducted at a stroke speed of 100 mm/sec and n = 3 for each bend radius. The presence of cracks in the outside portion of the bend part was determined with a stereoscopic microscope. The minimum bend radius at which cracking did not occur was considered the critical bend radius R. In the present invention, an evaluation of "Good" was given in terms of the bendability of the steel sheet in the case where the critical bend radius at 90°V bending R/t ≤ 2.5 (t: thickness of the steel sheet) was satisfied.

The high strength steel sheets prepared in Invention examples all had a TS of 980 MPa or more. In Invention examples, high strength steel sheets having excellent formability were prepared. In contrast, in Comparative examples, at least one of the properties TS, EL, λ, and bendability was poor.

[Table 4]

No.	Steel type	Thickness (mm)	Area fraction of F (%)	Area fraction of TM (%)	Area fraction of RA (%)	Area fraction of RA having aspect ratio of 2.0 or more and minor axis length of 1 µm or less/Total area fraction of RA	Average Mn content in RA (mass%)	Average Mn content in F (mass%)	Average Mn content in RA/Average Mn content in F	Average aspect ratio of RA	Average Mn content in RA/Average Mn content in F x Average aspect ratio of RA
1	A	12	55.6	18.9	20.6	0.341	6.57	2.12	3.10	5.41	16.77
2	A	1.4	53.5	20.0	19.0	0.390	6.80	2.85	2.39	4.35	10.38
3	A	1.6	54.4	20.4	22.2	0.492	6.47	2.78	2.33	4.89	11.38
4	A	12	48.8	24.4	17.6	0.546	6.30	2.73	2.31	5.12	11.82
5	A	12	49.8	18.0	22.3	0.175	6.55	2.18	3.00	4.44	13.34
6	A	1.4	28.8	15.4	21.8	0.413	529	3.14	1.68	3.99	6.72
7	A	1.4	45.8	37.3	16.3	0.304	6.77	2.80	2.42	4.55	11.00
8	A	1.6	45.9	30.3	15.2	0.195	6.12	2.45	250	1.03	2.57
9	A	1.4	51.2	25.3	15.9	0.215	5.82	2.91	2.00	122	24.4
10	A	12	54.2	19.2	20.2	0.497	8.77	2.61	3.36	3.69	12.39
11	A	1.4	55.3	15.4	195	0.181	6.32	28.7	220	124	2.73
12	A	12	55.9	15.5	19.6	0.137	6.10	2.82	2.16	1.35	2.92
13	A	1.4	54.9	18.9	20.3	0.592	723	1.72	4.20	5.14	21.61
14	B	1.4	50.8	20.3	22.1	0.613	5.05	123	411	421	17.28
15	C	1.6	51.6	18.5	24.9	0.444	6.10	1.82	3.35	3.55	1190
16	A	1.8	55.7	23.7	18.1	0.224	5.61	2.46	2.28	1.12	2.55
17	A	12	60.6	20.0	15.5	0.213	6.01	0.86	6.97	2.64	18.39
18	A	1.4	48.3	27.2	19.9	0.031	6.47	2.97	2.18	1.93	4.20
19	A	1.4	78.4	2.4	7.2	0.338	4.54	2.55	1.78	2.52	4.49
20	A	12	53.0	37.4	5.6	0.382	3.86	2.99	129	3.86	4.98
21	A	1.0	75.6	2.3	6.5	0.347	4.92	3.05	1.61	4.70	7.58
22	A	12	56.4	35.8	5.5	0.338	3.57	2.79	1.28	5.13	6.56
23	A	1.4	51.9	2.5	18.2	0.422	5.99	2.87	2.09	4.32	9.02
24	A	1.4	56.7	30.5	7.1	0.501	6.01	2.76	2.18	4.89	1065
25	A	12	45.3	2.1	17.7	0.409	6.11	2.80	2.18	5.62	12.26
26	D	1.4	60.5	18.9	18.4	0.630	5.38	2.71	1.99	5.38	10.68
27	E	1.4	53.2	14.6	23.4	0.484	4.54	2.78	1.64	6.20	1014
28	F	1.6	54.7	22.2	19.2	0.418	4.16	2.11	1.97	5.14	10.13
29	G	12	46.4	29.0	15.5	0.360	6.78	2.45	2.77	3.86	10.74
30	H	1.4	55.4	20.6	12.5	0.497	8.03	3.10	2.59	3.98	10.31
31	I	1.4	53.2	24.5	15.3	0.701	5.87	2.89	2.03	4.57	9.28
32	J	12	53.3	18.8	25.8	0.360	4.66	2.25	207	5.38	11.15
33	K	18	49.4	22.0	24.3	0.741	7.00	4.45	1.57	4.67	7.35
34	L	1.4	58.7	20.4	15.2	0.304	6.26	2.88	2.17	6.33	13.76
35	M	12	51.1	21.2	23.3	0.574	8.26	2.47	3.34	5.47	18.27
36	N	12	53.6	19.4	17.1	0.401	5.98	2.91	2.05	8.30	17.06
37	O	1.4	52.9	19.9	19.7	0.331	7.79	3.21	2.43	6.44	15.63
38	P	12	50.2	18.9	25.8	0.411	7.89	1.91	4.14	3.30	13.65
39	Q	1.4	48.5	28.2	14.2	0.352	4.84	1.07	4.52	4.98	22.51
40	R	1.4	79.8	2.5	7.4	0.316	6.01	2.44	2.46	2.89	7.12
41	S	1.4	51.5	14.8	25.0	0.667	7.16	2.67	2.68	5.12	13.73
42	S	12	53.1	14.4	25.9	0.646	6.86	2.97	2.31	4.46	10.30
43	T	12	49.7	15.3	26.3	0.306	9.95	6.12	1.63	6.49	1055
44	U	12	58.7	28.7	6.4	0.410	3.75	2.64	1.42	422	5.99
45	V	1.6	44.5	37.5	16.1	0.472	4.93	2.12	2.32	5.43	12.61
46	W	1.4	55.4	16.9	26.7	0.585	6.32	3.10	204	4.66	9.50
47	X	1.4	56.6	19.1	19.5	0.301	4.72	2.99	1.58	5.32	8.40
48	Y	1.6	52.7	19.8	24.5	0.305	5.24	2.41	2.17	429	9.33
49	Z	1.8	58.1	15.7	20.1	0.396	5.12	3.02	1.70	5.13	8.70
50	AA	12	51.0	22.0	23.6	0.409	5.25	2.55	2.06	6.19	12.75
51	AB	1.4	49.6	13.1	29.3	0.459	8.49	2.93	2.90	6.08	17.62
52	AC	1.4	46.3	30.1	20.3	0.489	7.29	3.03	2.41	5.44	13.09
53	AD	1.4	50.3	21.6	22.1	0.401	7.94	421	1.89	4.45	8.39
54	AE	12	54.1	17.6	22.8	0.378	6.15	2.06	2.96	5.34	15.79
55	AF	12	54.0	17.4	22.8	0.583	5.73	1.93	2.97	8.03	17.90
56	AG	1.4	52.2	21.4	21.3	0.422	4.96	3.04	1.63	5.90	9.63
57	AH	1.0	49.2	32.9	13.7	0582	5.08	2.95	1.72	4.07	7.01
58	AI	1.4	48.5	30.0	15.9	0.354	4.36	2.76	1.58	5.12	8.09
59	AJ	1.4	51.0	20.2	20.5	0.447	5.58	3.32	18.8	4.80	8.07
60	AK	1.4	49.5	20.4	20.2	0.328	6.41	327	1.96	5.40	10.59
61	AL	12	50.2	20.1	20.6	0311	6.87	2.60	2.64	4.13	10.91
62	AM	1.6	53.2	20.0	24.3	0.385	5.41	3.06	1.77	5.54	9.79
63	AN	1.4	52.7	20.3	20.6	0.335	4.77	1.69	2.82	4.52	12.76
64	AO	12	51.9	20.4	20.9	0.564	8.03	4.51	1.78	5.02	8.94

Underlined part Outside of the scope of the present invention
F: Ferrite, TM: Tempered martensite, RA: Retained austenite

[Table 5]

No.	Average C content in RA (mass%)	Average C content in F (mass%)	Average C content in RA/Average C content in F	Average C content in RA/Average Mn content in RA	Balance microstructure	TS (MPa)	EL (%)	λ (%)	R (mm)	R/t	Content of diffusive hydrogen in steel (mass ppm)	Remark
1	0.47	0.04	11.50	0.07	M,BF,P,θ	997	22.4	27	2.6		0.03	Invention example
2	0.39	0.04	9.75	0.06	M,BF,P,θ	1023	25.0	26	3.4	2.4	0.06	Invention example
3	0.40	0.05	8.35	0.06	M,BF,P,θ	1006	25.8	27	3.0	1.9	0.16	Invention example
4	0.44	0.12	3.67	0.07	M,BF,P,θ	1099	21.0	24	2.2	1.8	0.13	Invention example
5	0.43	0.06	7.17	0.07	M,BF,P,θ	1024	26.1	15	2.4	2.0	0.14	Comparative example
6	0.33	0.03	11.00	0.06	M,BF,P,θ	1055	14.2	32	3.0	2.1	0.03	Comparative example
7	0.34	0.09	3.78	0.05	M,BF,P,θ	1183	13.1	23	5.6	4.0	0.05	Comparative example
8	0.37	0.06	6.17	0.06	M,BF,P,θ	1175	14.4	11	2.2	1.4	0.08	Comparative example
9	0.36	0.06	6.29	0.06	M,BF,P,θ	1172	13.2	14	3.2	2.3	0.02	Comparative example
10	0.19	0.04	4.54	0.02	M,BF,P,θ	1021	23.8	26	5.0	4.2	0.04	Comparative example
11	0.40	0.03	14.62	0.08	M,BF,P,θ	1052	18.5	19	32	2.3	0.10	Comparative example
12	0.46	0.08	5.98	0.08	M,BF,P,θ	983	18.9	19	1.6	1.3	0.15	Comparative example
13	0.43	0.09	4.78	0.06	M,BF,P,θ	1030	23.7	29	1.6	1.1	0.12	Invention example
14	0.44	0.10	4.40	0.09	M,BF,P,θ	1050	22.5	28	2.8	2.0	0.23	Invention example
15	0.45	0.09	5.00	0.07	M,BF,P,θ	1016	27.3	30	2.6	1.6	0.11	Invention example
16	0.46	0.06	7.51	0.08	M,BF,P,θ	1131	15.3	11	2.6	1.4	0.20	Comparative example
17	0.38	0.09	4.22	0.06	M,BF,P,θ	1021	24.6	12	1.8	1.5	0.18	Comparative example
18	0.48	0.12	4.00	0.07	M,BF,P,θ	985	25.7	16	3.2	2.3	0.15	Comparative example
19	0.33	0.07	4.71	0.07	M,BF,P,θ	943	29.7	33	0.8	0.6	0.09	Comparative example
20	0.33	0.14	2.36	0.09	M,BF,P,θ	1155	12.6	19	4.0	3.3	0.12	Comparative example
21	0.35	0.11	3.18	0.06	M,BF,P,θ	902	25.6	35	1.0	1.0	0.22	Comparative example
22	0.27	0.12	2.25	0.08	M,BF,P,θ	1123	13.7	24	3.8	3.2	0.06	Comparative example
23	0.45	0.06	7.50	0.08	M,BF,P,θ	1004	24.5	14	4.5	3.2	0.58	Comparative example
24	0.41	0.05	8.20	0.07	M,BF,P,θ	1022	16.4	26	2.0	1.4	0.04	Comparative example
25	0.41	0.08	5.13	0.07	M,BF,P,θ	1039	21.2	16	4.0	3.3	0.51	Comparative example
26	0.30	0.07	4.01	0.06	M,BF,P,θ	1046	20.7	22	2.4	1.7	0.23	Invention example
27	0.37	0.04	9.25	0.08	M,BF,P,θ	982	22.5	30	3.2	2.3	0.24	Invention example
28	0.38	0.09	4.22	0.09	M,BF,P,θ	1160	22.2	24	1.2	0.8	0.20	Invention example
29	0.50	0.06	10.00	0.07	M,BF,P,θ	1078	23.7	24	2.6	2.1	0.24	Invention example
30	0.46	0.04	11.50	0.06	M,BF,P,θ	1000	20.9	31	0.6	0.5	0.23	Invention example
31	0.31 .	0.07	4.43 6	0.05	M,BF,P,θ	1035	24.1	23	2.8	2.0	0.09	Invention example
32	0.40	0.09	5.46	0.10	M,BF,P,θ	988	26.2	23	2.6	2.2	0.17	Invention example
33	0.38	0.08	4.75	0.05	M,BF,P,θ	1003	28.0	28	3.2	1.8	0.02	Invention example
34	0.47	0.05	9.40	0.08	M,BF,P,θ	1085	21.0	25	2.3	1.7	0.05	Invention example
35	0.46	0.11	4.11	0.06	M,BF,P,θ	1052	24.1	27	2.2	1.8	0.14	Invention example
36	0.53	0.04	13.25	0.09	M,BF,P,θ	1021	22.1	26	2.8	2.3	0.12	Invention example
37	0.41	0.07	5.86	0.05	M,BF,P,θ	995	21.6	25	3.0	2.1	0.29	Invention example
38	0.44	0.05	8.80	0.06	M,BF,P,θ	1024	22.9	25	2.2	1.8	0.12	Invention example
39	0.25	0.06	4.17	0.05	M,BF,P,θ	1064	21.4	23	2.6	1.9	0.18	Invention example
40	0.33	0.02	16.50	0.06	M,BF,P,θ	841	15.2	26	2.2	1.6	0.05	Comparative example
41	0.37	0.04	9.23	0.05	M,BF,P,θ	989	15.1	32	5.0	3.6	0.21	Comparative example
42	0.47	0.06	7.83	0.07	M,BF,P,θ	984	16.0	26	4.0	3.3	0.22	Comparative example
43	0.37	0.07	5.29	0.04	M,BF,P,θ	1002	14.8	14	3.5	2.9	0.10	Comparative example
44	0.42	0.06	7.00	0.11	M,BF,P,θ	1025	13.5	20	2.6	2.2	0.10	Comparative example
45	0.52	0.11	4.73	0.11	M,BF,P,θ	1023	24.5	13	2.4	1.5	0.24	Comparative example
46	0.44	0.05	8.80	0.07	M,BF,P,θ	1012	25.9	24	1.8	1.3	0.08	Invention example
47	0.49	0.05	9.90	0.10	M,BF,P,θ	1129	23.2	28	2.4	1.7	0.19	Invention example
48	0.51	0.06	8.50	0.10	M,BF,P,θ	1000	22.9	25	2.6	1.6	0.23	Invention example
49	0.51	0.07	7.29	0.10	M,BF,P,θ	988	21.3	27	2.6	1.4	0.28	Invention example
50	0.54	0.07	7.71	0.10	M,BF,P,θ	1075	23.2	30	2.2	1.8	0.19	Invention example
51	0.44	0.09	4.89	0.05	M,BF,P,θ	983	30.4	25	2.6	1.9	0.09	Invention example
52	0.42	0.04	10.50	0.06	M,BF,P,θ	1082	20.3	18	2.6	1.9	0.03	Invention example
53	0.42	0.04	10.50	0.05	M,BF,P,θ	1014	21.0	25	1.2	0.9	0.22	Invention example
54	0.50	0.04	12.50	0.08	M,BF,P,θ	1002	21.4	28	2.6	2.2	0.15	Invention example
55	0.33	0.06	5.32	0.06	M,BF,P,θ	997	24.3	29	1.6	1.3	0.22	Invention example
56	0.45	0.09	5.00	0.09	M,BF,P,θ	1104	24.5)	27	1.2	0.9	0.16	Invention example
57	0.34	0.10	3.50	0.07	M,BF,P,θ	1089	21.2	25	2.0	2.0	0.09	Invention example
58	0.41	0.06	6.83	0.09	M,BF,P,θ	1033	22.4	24	2.4	1.7	0.28	Invention example
59	0.37	0.12	3.08	0.07	M,BF,P,θ	997	26.7	27	2.8	2.0	0.11	Invention example
60	0.43	0.10	4.14	0.07	M,BF,P,θ	982	23.1	28	2.4	1.7	0.09	Invention example
61	0.46	0.07	6.57	0.07	M,BF,P,θ	1028	22	29	2.2	1.8	0.08	Invention example
62	0.45	0.09	5.00	0.08	M,BF,P,θ	995	26.8	27	2.8	1.8	0.06	Invention example
63	0.40	0.07	5.71	0.08	M,BF,P,θ	991	20.9	29	2.2	1.6	0.07	Invention example
64	0.44	0.03	14.67	0.06	M,BF,P,θ	986	22.1	28	2.8	2.3	0.08	Invention example

Underlined part: Outside of the scope of the present invention
F: Ferrite, M: Martensite, RA: Retained austenite, BF: Bainitic ferrite, P: Pearlite, θ: Carbides (e.g., cementite)

Industrial Applicability

According to the present invention, a high strength steel sheet having a TS (tensile strength) of 980 MPa or more and excellent formability can be produced. Applying the high strength steel sheet according to the present invention to, for example, automotive structural components reduces the weights of car bodies and thereby improves fuel economy. Thus, the use of the high strength steel sheet according to the present invention is highly valuable from an industrial viewpoint.

Claims

A high strength steel sheet comprising:
a chemical composition containing, by mass,

C: 0.030% or more and 0.250% or less,

Si: 0.01% or more and 3.00% or less,

Mn: 2.50% or more and 8.00% or less,

P: 0.001% or more and 0.100% or less,

S: 0.0001% or more and 0.0200% or less,

N: 0.0005% or more and 0.0100% or less, and

Al: 0.001% or more and 2.000% or less,

with the balance being Fe and incidental impurities;

a steel microstructure including, by area, ferrite: 30% or more and 80% or less, tempered martensite: 3.0% or more and 35% or less, and retained austenite: 8% or more,

wherein a quotient of an area fraction of grains of the retained austenite, the grains having an aspect ratio of 2.0 or more and a minor axis length of 1 µm or less, divided by a total area fraction of the retained austenite is 0.3 or more,

wherein a quotient of an average Mn content (mass%) in the retained austenite divided by an average Mn content (mass%) in the ferrite is 1.5 or more, and a product of the quotient of the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite and an average aspect ratio of the retained austenite is 3.0 or more,

wherein a quotient of an average C content (mass%) in the retained austenite divided by an average C content (mass%) in the ferrite is 3.0 or more,

wherein a quotient of the average C content (mass%) in the retained austenite divided by the average Mn content (mass%) in the retained austenite is 0.05 or more, and

wherein a content of diffusible hydrogen in steel is 0.3 ppm by mass or less.
The high strength steel sheet according to Claim 1, wherein the chemical composition further contains at least one element selected from, by mass,
Ti: 0.200% or less,

Nb: 0.200% or less,

V: 0.500% or less,

W: 0.500% or less,

B: 0.0050% or less,

Ni: 1.000% or less,

Cr: 1.000% or less,

Mo: 1.000% or less,

Cu: 1.000% or less,

Sn: 0.200% or less,

Sb: 0.200% or less,

Ta: 0.100% or less,

Ca: 0.0050% or less,

Mg: 0.0050% or less,

Zr: 0.0050% or less, and

REM: 0.0050% or less.
The high strength steel sheet according to Claim 1 or 2, the high strength steel sheet further comprising a galvanized layer disposed on a surface of the high strength steel sheet.
The high strength steel sheet according to Claim 3, wherein the galvanized layer is a galvannealed layer.
A method for manufacturing a high strength steel sheet, the method comprising heating a steel slab having the chemical composition according to Claim 1 or 2, hot rolling the steel slab at a finish rolling delivery temperature of 750°C or more and 1000°C or less, then performing coiling at 300°C or more and 750°C or less, subsequently performing cold rolling, then performing holding at a temperature equal to or higher than an Ac₃ transformation temperature - 50°C for 20 s or more and 1800 s or less, subsequently performing cooling to a cooling stop temperature equal to or lower than a martensitic transformation start temperature, then performing reheating to a reheating temperature of 120°C or more and 450°C or less, subsequently performing holding at the reheating temperature for 2 s or more and 600 s or less, then performing cooling to room temperature, subsequently performing holding at a temperature equal to or lower than an Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less, then performing cooling to 100°C or less, then performing holding at higher than 100°C and 400°C or less for 10 s or more, and subsequently performing cooling.
The method for manufacturing a high strength steel sheet according to Claim 5, the method further comprising, subsequent to the coiling and prior to the cold rolling, performing holding at a temperature equal to or lower than the Ac₁ transformation temperature for more than 1800 s.
The method for manufacturing a high strength steel sheet according to Claim 5 or 6, the method further comprising, subsequent to the holding at the temperature equal to or higher than the Ac₁ transformation temperature and equal to or lower than the Ac₁ transformation temperature + 150°C for 20 s or more and 600 s or less, performing cooling, then performing a galvanizing treatment, and subsequently performing the cooling to 100°C or less.
The method for manufacturing a high strength steel sheet according to Claim 7, the method further comprising, subsequent to the galvanizing treatment, performing an alloying treatment at 450°C or more and 600°C or less.